KR101819358B1 - High-strength thin steel sheet having excellent formability and method for manufacturing the same - Google Patents

Abstract

The present invention relates to a high-strength thin steel sheet having excellent formability and a manufacturing method thereof. The high-strength thin steel sheet comprises 0.001-0.004 wt% of C, 0.5 wt% or lower (excluding 0 wt%) of Si, 1.2 wt% or lower (excluding 0 wt%) of Mn, 0.005-0.12 wt% of P, 0.01 wt% or lower of S, 0.01 wt% or lower of N, 0.1 wt% or lower (excluding 0 wt%) of acid soluble Al, 0.01-0.04 wt% of Ti, and the remainder consisting of Fe and inevitable impurities. A content of Ti, N and S satisfies following interaction formula 1. A ratio (b/a) of an average random strength ratio (b) of a (111)[1-10]~(111)[-1-12] orientation group to an average random strength ratio (a) of a (001)[1-10]~(110)[1-10] orientation group at a t/4 (t is a thickness of the thin steel sheet) point in a sheet thickness direction is 2.3 or higher. Bake hardenability (BH) is 4 MPa or higher. [Interaction formula 1] -0.02<=[Ti]-(24/7)[N]-(3/2)[S]<=0.025 (each of [Ti], [N] and [S] is means a content of the corresponding element in wt%).

Description

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a high strength steel sheet having excellent moldability and a method of manufacturing the steel sheet. 2. Description of the Related Art HIGH- STRENGTH THIN STEEL SHEET HAVING EXCELLENT FORMABILITY AND METHOD FOR MANUFACTURING THE SAME

The present invention relates to a high-strength thin-gauge steel sheet and a method of manufacturing the same, and more particularly, to a high-strength thin-gauge steel sheet excellent in moldability which can be preferably applied to a material for automobile outer plates and the like, and a method of manufacturing the same.

Steel that is used as a material for inner and outer plates (doors, hoods, fenders, floors, etc.) of automobiles is required not only high strength but also good formability. This is to ensure the safety of passengers from accidents, and to improve the fuel efficiency by lightening the weight of the vehicle.

However, it is very difficult to satisfy both of the above factors (strength and formability) at the same time because the increase in the strength of the steel sheet causes deterioration of the formability. Particularly, Cracks occur during machining in the parts requiring the formation, and therefore the application of high strength steels to these parts is still insufficient.

A known steel sheet having excellent strength and formability developed so far includes so-called Interstitial Free Steel. This is achieved by adding strong carbonitride forming elements such as titanium (Ti) and / or niobium (Nb) to remove solid elements such as carbon (C), nitrogen (N) and sulfur (S) And are typically disclosed in Patent Documents 1 to 4. However, the IF steel exhibits an average plastic anisotropy coefficient (rankford value, r value) of 1.5 to 1.8, which is insufficient to substitute for a conventional cold rolled steel sheet having a deep drawing quality (DDQ) grade.

Japanese Patent Application Laid-Open No. 1992-280943 Japanese Laid-Open Patent Publication No. 1993-070836 Japanese Patent Application Laid-Open No. 1993-263184 Japanese Laid-Open Patent Publication No. 1998-096051

One of the objects of the present invention is to provide a high strength thin steel sheet excellent in moldability and a method of manufacturing the same.

An aspect of the present invention is a steel sheet comprising, by weight%, 0.001 to 0.004% of C, 0.5% or less of Si (excluding 0%), 1.2% or less of Mn (excluding 0%), 0.005 to 0.12% % Of Ti, 0.01 to 0.04% of Ti, and the balance of Fe and inevitable impurities, and the contents of Ti, N and S satisfy the following relational expression 1 And the average random intensity ratio (a) of the (001) [1-10] to (110) [1-10] bearing group at t / 4 (t: thickness of thin steel plate) 111) A high strength thin steel plate having a ratio (b / a) of an average random intensity ratio (b) of [1-10] to (111) [1-12] bearing groups of 2.3 or more and a bake hardening amount .

[Relation 1] -0.02? [Ti] - (24/7) [N] - (3/2) [S]? 0.025

(Where each of [Ti], [N] and [S] means the content (weight%) of the corresponding element)

Another aspect of the present invention provides a method of manufacturing a semiconductor device, comprising: 0.001 to 0.004% of C, 0.5% or less of Si (excluding 0%), 1.2% or less of Mn (excluding 0%), 0.005 to 0.12% Hot rolling a steel slab containing 0.01% or less of N, 0.01% or less of N, 0.1% or less of acid soluble Al (excluding 0%), 0.01 to 0.04% of Ti, Fe and unavoidable impurities, Rolling the rolled hot-rolled steel sheet at a temperature of 450 to 750 ° C, cold-rolling the rolled hot-rolled steel sheet at a reduction ratio of 75% or more to obtain a cold-rolled steel sheet, raising the cold-rolled steel sheet to an annealing temperature of 830 to 880 ° C Continuously cooling the cold-rolled steel sheet at a temperature of 650 ° C or lower at a rate of 2 to 10 ° C / sec, cooling the cold-rolled steel sheet to a temperature of 650 ° C or lower, And temper rolling the steel sheet at a reduction ratio of 0.3 to 1.6%, wherein at the time of heating the cold-rolled steel sheet, from the (recrystallization starting temperature + 20) The average cooling rate to provide a method for manufacturing a high strength steel sheet characterized in that not more than 5 ℃ / sec.

As one of the various effects of the present invention, the thin steel sheet according to the present invention is excellent in strength and moldability and can be suitably applied to materials for automobile exterior panels and the like.

FIG. 1 is a graph showing the developmental degree of the group organization of Inventive Example 1. FIG.

As a result of deep research to solve the problems of the prior art described above, the inventors of the present invention have found that the addition of titanium (Ti), which is a strong carbonitride forming element in steel, alone or in combination with titanium (Ti) and niobium It is necessary to remove the solid elements such as carbon (C), nitrogen (N) and sulfur (S), and appropriately control the position distribution of the carbide or the like generated as a result of the removal of the solid element, It is possible to remarkably improve the hardenability of the sintered body, and it has been confirmed that the hardening of the sintering can be remarkably improved by leaving the dissolved carbon to be remained at an appropriate level during the annealing, and the present invention has been accomplished.

BEST MODE FOR CARRYING OUT THE INVENTION Hereinafter, a high strength steel sheet excellent in formability which is one aspect of the present invention will be described in detail.

First, the alloy component and the preferable content range of the high strength steel sheet will be described in detail. It is to be noted that the content of each component described below is based on weight unless otherwise specified.

C: 0.001 to 0.004%

Carbon is an invasive solid element and has a great influence on the formation of aggregate structure of steel sheet during cold rolling and annealing. Particularly, when the amount of solid carbon in the steel increases, the growth of grains having a {111} texture that is advantageous for drawing workability is suppressed and the growth of grains having {110} and {100} texture is promoted, Can be deteriorated. Further, when the carbon content is excessively excessive, the Ti content necessary for precipitating the carbide is excessively disadvantageous from the viewpoint of economical efficiency, and there is also a problem that the fine TiC carbide is distributed in a large amount in the steel and drastically deteriorates the drawability. Therefore, in the present invention, the upper limit of the carbon content is controlled to 0.004%, preferably 0.0035%. On the other hand, the lower the carbon content, the better the improvement of the drawability, but if the content is too low, the hardening of the hardening of the thin steel sheet drastically deteriorates. Therefore, in the present invention, the lower limit of the carbon content is controlled to 0.001%, preferably 0.0012%.

Si: 0.5% or less (excluding 0%)

Silicon contributes to the strength increase of the steel sheet by solid solution strengthening. However, if the content thereof is excessive, there is a problem that surface scale defects are caused to deteriorate the surface characteristics of the plating. Therefore, in the present invention, the upper limit is controlled to 0.5%, preferably 0.05%. In the present invention, the lower limit of the silicon content is not particularly limited, but may be 0.001%, and more preferably 0.002%.

Mn: 1.2% or less (excluding 0%)

Manganese is a solid solution strengthening element which not only contributes to the improvement of steel strength but also precipitates S in MnS to inhibit plate breakage and high temperature embrittlement by S during hot rolling. However, when the content thereof is excessive, excess Mn is solved to deteriorate drawability. In the present invention, the upper limit of the manganese content is controlled to 1.2% or less, preferably 1.0% or less, and more preferably 0.8% or less. In the present invention, the lower limit of the manganese content is not particularly limited, but may be preferably 0.01%, more preferably 0.1%.

P: 0.005 to 0.12%

Phosphorus is the most effective element to improve the strength of the steel, with a very good employment effect and without significantly deteriorating the drawability. In the present invention, the lower limit of the phosphorus content is controlled to 0.005%, preferably 0.008%, and more preferably 0.010%. However, when the content is excessive, there is a problem that excessive P precipitates in FeTiP and drawability deteriorates. In the present invention, the upper limit of phosphorus content is controlled to 0.12%, preferably to 0.10%, and more preferably to 0.08%.

S: 0.01% or less, N: 0.01% or less

Sulfur and nitrogen are inevitably impurities in the steel. In order to secure excellent welding characteristics, it is desirable to control their contents as low as possible. In the present invention, the upper limit of the content of sulfur and nitrogen is controlled to 0.01% or less, respectively, in terms of ensuring proper welding characteristics.

Sol.Al: 0.1% or less (excluding 0%)

Acid soluble aluminum precipitates AlN and contributes to improving the drawability and ductility of the thin steel sheet. However, when the content is excessive, Al-based inclusions are excessively formed at the time of steelmaking, thereby causing defects in the steel sheet. In the present invention, the upper limit of the content of acid soluble aluminum is controlled to 0.1%, preferably 0.08%, and more preferably 0.05%. In the present invention, the lower limit of the content of the acid soluble aluminum is not particularly limited, but may be preferably 0.01%, more preferably 0.02%.

Ti: 0.01 to 0.04%

Titanium is an element contributing to improving the drawability of the thin steel sheet by precipitating the Ti-based carbonitrides by reacting with the solid carbon and the solid nitrogen in the hot-rolled steel. In the present invention, the lower limit of the titanium content is controlled to 0.01% or more, preferably 0.012% or more, and more preferably 0.015% or more. However, there is a fear that the TiC or TiN precipitates are distributed in a large amount in the steel due to the excessive amount of the TiC or TiN precipitates. The amount of carbon is excessively low, and there is a fear that the hardening of the bare steel sheet may be deteriorated. In the present invention, the upper limit of the titanium content is controlled to 0.04%, preferably 0.03%.

And the balance Fe and inevitable impurities. However, in the ordinary manufacturing process, impurities which are not intended from the raw material or the surrounding environment may be inevitably incorporated, so that it can not be excluded. These impurities are not specifically referred to in this specification, as they are known to one of ordinary skill in the art. In addition, addition of an effective component other than the above-mentioned composition is not excluded, and in particular, the following components may be further included to further improve the mechanical properties of the steel sheet.

Nb: 0.005 to 0.04%

Niobium functions to facilitate the formation of aggregate structure during annealing by precipitating solid carbon in the form of (Ti, Nb) C complex carbide in hot rolling. Further, when an appropriate amount of Nb is added, plastic deformation anisotropy (0 °, 45 °, 90 °) is improved, and plastic deformation anisotropy (r-value) in the 0 ° and 45 ° directions with respect to the 90 ° direction is improved. As a result, the planar anisotropy of the material reaches approximately zero, and the r value is evenly distributed on the surface of the plate, so that the molding of the ear shape of the material at the time of molding Defects can be prevented. In order to obtain such an effect in the present invention, the lower limit of the niobium content is preferably controlled to 0.005% or more, more preferably 0.008% or more. However, when the content is excessively excessive, most of the solid carbon in the steel is precipitated as fine NbC, so that the solidified carbon is hardly re-dissolved even after the annealing, so that the hardening of the hardening is deteriorated. Further, the precipitation amount of fine (Ti, Nb) There is a problem that not only the drawability (r-value) is deteriorated but also the material deterioration due to the increase in the recrystallization temperature is caused. The upper limit of the niobium content is preferably 0.04%, more preferably 0.03%, even more preferably 0.025%.

B: 0.002% or less (excluding 0%)

Boron suppresses secondary machining brittleness due to P in the steel. However, if the content is excessive, the ductility of the steel sheet may be lowered. In the present invention, the upper limit of the boron content is controlled to 0.002% or less, preferably 0.0015% or less. In the present invention, the lower limit of the boron content is not particularly limited, but may be 0.0003%, and more preferably 0.0005%.

On the other hand, at the time of designing the alloy of the thin steel sheet having the above-mentioned composition range, it is preferable that the contents of Ti, N and S satisfy the following relational expression (1). If the value of [Ti] - (24/7) [N] - (3/2) [S] is less than -0.02, the Ti content for precipitation of steel C to TiC is absolutely insufficient, On the other hand, when the value exceeds 0.025, FeTiP precipitates are formed in addition to the TiC precipitates favorable in workability, and the development of the {111} orientation is significantly inhibited during annealing. More preferably, the value is controlled to be -0.01 to 0.01.

[Relation 1] -0.02? [Ti] - (24/7) [N] - (3/2) [S]? 0.025

(Where each of [Ti], [N] and [S] means the content (weight%) of the corresponding element)

Hereinafter, the structure and the precipitate of the high-strength steel sheet will be described in detail.

An array having a certain plane and orientation generated inside a crystal is called a texture, and a pattern in which these aggregate structures develop into a band in a certain direction is called a fiber aggregate structure. The texture of the texture is closely related to the drawability and it is known that the higher the surface strength value of the gamma (γ) -fiber aggregate structure in which {111} faces are formed perpendicular to the rolling face, the better the drawing processability is . The alpha (alpha) -fiber assembly texture is usually defined as RD // <110> and the gamma (γ) -fiber assembly texture is defined as ND // <111>.

On the other hand, in order to form gamma-fiber aggregate structure as described above, the present inventors have found that alpha (alpha) -fibrous texture ((001 (Γ) -fiber aggregate structure ((111) [1-10] to (111) [- 1] of the average random intensity ratio (a) -12] defense group) is very important. More specifically, the average random intensity ratio (a) of the (001) [1-10] to (110) [1-10] bearing group at a point of t / 4 (B / a) of the average random intensity ratio (b / a) of the (111) [1-10] to (111) [- 1-12] bearing group to the average random anisotropy coefficient , r value) of 1.9 or more can be ensured and it is confirmed that excellent drawability can be secured. On the other hand, the higher the average random intensity ratio of gamma-fiber aggregate ((111) [1-10] to (111) [1-12] bearing group) And is not particularly limited.

Particularly, in the present invention, it is confirmed that, when molding an automobile part, it is necessary to ensure excellent drawability in various directions, not in a specific direction, it is possible to form a complete part without cracking, and the development degree of gamma- It was found that the complete moldability can be expressed by displaying the value. (111) [1-10], 30 ((111) [1-21]), 60 ((111) [0-11]), 90 ((111) [- 1-12]), the more the development of the average random intensity ratio is generally higher.

On the other hand, the average plastic anisotropy coefficient (Lankford value, r value) obtained from the plastic anisotropy coefficient measured for each direction with respect to the rolling direction is a representative material characteristic value indicating drawability.

r value = (r0 + r90 + 2r45) / 4 (1)

(Where ri represents the plastic anisotropy coefficient measured in the specimen taken in the direction of i from the rolling direction).

The larger the value of r in the above equation, the greater the depth of the molding cup in the drawing process, and the better the drawability can be judged. The thin steel sheet according to one embodiment of the present invention has an r value of 1.9 or more and exhibits excellent drawability.

According to one example, the average grain size of the high-strength thin steel sheet may be 5 탆 or more, and preferably 7 탆 or more. Here, the average grain size means an equivalent circular diameter of the crystal grains. In the present invention, it is advantageous to ensure coarser crystal grains as much as possible as the grain size becomes coarser in terms of moldability. For this purpose, the C content is reduced to an extremely low carbon steel level of 40 ppm or less through component control, and the carbide precipitation is controlled as effectively as possible to achieve crystal growth during annealing. This is because the larger the grain size is, the easier the precipitation of carbide in the crystal grain than the grain boundary system, and the possibility of cracking during processing can be remarkably lowered. On the other hand, the larger the average crystal grain size is, the more advantageous from the moldability standpoint. In the present invention, the upper limit of the average crystal grain size is not particularly limited, but there is a fear of damage to the refractory bricks in the annealing furnace due to high- Considering the side, the upper limit can be limited to 20 탆.

According to one example, the high strength steel sheet of the present invention may have a P value of 80% or more, preferably 82% or more, defined by the following formula (1). When the ratio (P) is less than 80%, that is, when a large amount of carbides are precipitated in the grain boundaries, the possibility of occurrence of cracks during machining becomes considerably high, which may deteriorate ductility and drawability. The higher the ratio P is, the better the ductility and the drawability are improved. Therefore, in the present invention, the upper limit of the ratio (P) is not particularly limited. On the other hand, the term "carbide" as used herein means TiC single carbide, NbC single carbide or (Ti, Nb) C complex carbide .

[Equation 1]

P _in (%) = {N _in / (N _in + N _gb )} x 100

(Where N _in is the number of carbides having a circle equivalent diameter of 20 nm or less and N _gb is the number of carbides having a circle equivalent diameter of 20 nm or less present in grain boundaries)

According to one example, the high strength steel sheet of the present invention may contain 0.2 or less FeTiP precipitates per unit area (탆 ² ), preferably 0.1 Or less. The FeTiP precipitates mainly precipitate in the form of needle, which deteriorates the development of {111} orientation during annealing. If the formation beyond this the precipitate FeTiP 0. 2 / ㎛ ^2, there is a fear that the drawability deteriorates. On the other hand, the smaller the number of FeTiP precipitates per unit area is, the better the drawability is improved. Therefore, in the present invention, the lower limit of the number of FeTiP precipitates is not particularly limited.

According to one example, the high-strength thin steel sheet of the present invention has a strength of 4 MPa or higher, (BH) of 10 MPa or more, and more preferably 15 MPa or more, and exhibits excellent bake hardenability.

According to one example, the high strength steel sheet of the present invention has a thickness of 0.8 mm or less and has a product of a yield strength (YS, MPa) and an average plastic anisotropy coefficient (Lankford value, r-value) of 290 MPa or more, Dent property and moldability, which means resistance to force, are very excellent and can be preferably applied to a material for an automobile shell plating.

The high strength steel sheet of the present invention described above can be manufactured by various methods, and the manufacturing method thereof is not particularly limited. However, as a preferable example, it can be produced by the following method.

Hereinafter, a method of manufacturing a high strength steel sheet excellent in formability, which is another aspect of the present invention, will be described in detail.

First, a hot-rolled steel sheet is obtained by hot-rolling a steel slab having the above-mentioned component system.

According to one example, the finish rolling during hot rolling can be carried out at austenite single phase reverse temperature (Ar3 or higher temperature). If the hot finish rolling temperature is less than Ar 3, there is a high possibility of rolling in two-sided rolling, which may cause material non-uniformity. For reference, Ar3 can be calculated from Equation 2 below.

[Formula 2]

15 [Cr] - 55 [Ni] - 80 [Mo] - 20 [Cu] - 15 [

(Where each of [C], [Mn], [Cu], [Cr]. [No]

Next, the hot-rolled steel sheet is wound.

At this time, the coiling temperature is preferably 450 to 750 占 폚, more preferably 500 to 700 占 폚. If the coiling temperature is less than 450 ° C, a large amount of FeTiP precipitates may be precipitated to lower the drawability and the plate warping may occur. On the other hand, when the coiling temperature exceeds 750 ° C, There is a possibility that the hardening amount BH of the hardening is deteriorated.

According to one example, the average cooling rate from the hot finish rolling temperature to the coiling temperature may be 10-200 占 폚 / sec. If the average cooling rate is less than 10 ° C / sec, the ferrite grains grow unevenly and FeTiP precipitates are formed, which makes it difficult to secure the desired formability of the present invention. On the other hand, when the cooling rate exceeds 200 ° C / The temperature of the hot-rolled steel sheet may become uneven and the shape of the hot-rolled steel sheet may become poor.

Next, the rolled hot-rolled steel sheet is cold-rolled to obtain a cold-rolled steel sheet.

At this time, the cold rolling reduction rate is preferably 75% or more. If the cold rolling reduction is less than 75%, there is a problem that the gamma (γ) -fiber aggregate structure does not sufficiently grow and the drawability becomes poor. On the other hand, the upper limit of the cold rolling reduction rate is not particularly limited in the present invention because the higher the cold rolling reduction rate is, the more favorable the gamma (?) - fiber aggregate texture growth. However, if the cold rolling reduction rate is too high, The shape may become poor, and when this is considered, the upper limit may be limited to 85%.

Next, the cold-rolled steel sheet is continuously annealed.

At this time, the annealing temperature (T) is preferably 830 to 880 ° C, and more preferably 840 to 870 ° C. If the annealing temperature (T) is less than 830 占 폚, the gamma (?) - fiber aggregate structure which is advantageous for workability may not grow sufficiently and the drawability may be weakened. The precipitate during annealing may not be redissolved, ) May be deteriorated. On the other hand, if the annealing temperature (T) exceeds 880 DEG C, it may be advantageous in workability, but the shape of the steel sheet becomes poor due to grain size deviation, and there is a possibility of causing a problem in the annealing heating furnace.

On the other hand, the annealing time (t), that is, the holding time at the annealing temperature is preferably 30 to 80 sec, more preferably 40 to 70 sec. When the annealing time is sufficiently secured after sufficiently developing the gamma (?) - fiber aggregate structure, some carbides are reused as the solid carbon. When cooling is performed in the state where such solid carbon exists, (BH) of the cured product. If the annealing time t is less than 30 sec, the amount of the cured carbon in the thin steel sheet does not remain or is insufficient due to the shortage of the redissolving time. On the other hand, if the annealing time t exceeds 80 sec, The crystal grains are coarsened and the crystal grain size deviation causes the steel sheet to be inferior in shape, which is disadvantageous in terms of economy.

According to one example, at the time of continuous annealing, the annealing temperature (T, 占 폚) and the annealing time (t, sec) can satisfy the following relational expression (2). If the value of 0.001 * T * t is less than 30, drawability and sintering hardenability may deteriorate. On the other hand, when the value of 0.001 * T * t exceeds 70, crystal grain coarsening and grain size deviation occur, .

[Relation 2] 30? 0.001 * T * t? 70

On the other hand, at the time of continuous annealing, the average heating rate from the recrystallization starting temperature + 20 占 폚 to the annealing temperature is preferably 5 占 폚 / sec or less, more preferably 4.5 占 폚 / sec or less and still more preferably 3.8 占 폚 / sec or less. Here, the recrystallization starting temperature is defined as a temperature at which a new recrystallized grain begins to be formed in the process of annealing a rolled structure elongated by cold rolling. More specifically, the area fraction of new recrystallized grains in the entire grain is 50% Is defined as the temperature at the point of occupancy. In the initial stage of recrystallization, nucleation and growth of new crystal grains are accompanied. The lower the rate of temperature rise at this stage, the more nucleation of {111} . If the rate of temperature rise in the above-mentioned temperature range exceeds 5 DEG C / sec, nucleation of the {111} texture is not sufficient at the time of recrystallization, and the crystal grains are refined to sufficiently secure the workability required in the present invention There is a fear of not being able to do. On the other hand, the slower the rate of temperature rise in the above-mentioned temperature range, the more advantageous nucleation and nucleation of the {111} texture that is advantageous to the workability is not particularly limited in the present invention.

Next, the cold-rolled steel sheet subjected to the continuous annealing is cooled to a temperature of 650 DEG C or lower.

At this time, the average cooling rate is preferably 2 to 10 ° C / sec, more preferably 3 to 8 ° C / sec. If the average cooling rate is less than 2 캜 / sec, there is a fear that the dissolved carbon that is redissolved during the annealing will be re-precipitated as carbide and the hardening property of the hardening may be deteriorated. On the other hand, There is a concern. On the other hand, 650 ° C is the temperature at which most of the carbide precipitation and diffusion is completed, and the cooling conditions thereafter are not particularly limited.

Next, the cooled cold-rolled steel sheet is temper rolled to obtain a high-strength thin-rolled steel sheet.

At this time, the rough reduction ratio is preferably 0.3 to 1.6%. The temper rolling increases the yield strength of the steel, increases the endurance by a large amount of the mobilized potential introduced during rolling, and increases the hardening of the sinter by the interaction of the solid carbon with the dislocations. If the rough reduction ratio is less than 0.3%, it is not only disadvantageous to the plate shape control but also a possibility that a stretch strain defect is likely to occur because the movable potential is not secured sufficiently. On the other hand, when the reduction ratio is more than 1.6% Not only the possibility increases but also the tendency to decrease to the r-value which is the formability index.

Next, a hot-dip galvanized steel sheet may be obtained by performing hot-dip galvanizing on the surface of the high-strength steel sheet, if necessary, or subjected to alloying heat treatment after hot-dip galvanizing, to obtain an alloyed hot-dip galvanized steel sheet. At this time, the alloying heat treatment temperature is preferably 450 to 600 ° C. If the alloying heat treatment temperature is less than 450 ° C, alloying is not sufficient, which may result in a sacrificial action or deterioration of the plating adhesion. On the other hand, when the temperature exceeds 600 ° C, alloying becomes excessive and degradation of powdering property . On the other hand, the Fe concentration in the plating layer after the alloying heat treatment is preferably 8 to 12 wt%.

Hereinafter, the present invention will be described in more detail with reference to Examples. However, the following examples are only illustrative of the present invention in more detail and do not limit the scope of the present invention.

( Example )

A steel slab (220 mm in thickness) having the alloy composition shown in the following Table 1 was heated to 1200 캜 and hot-rolled to produce a hot-rolled steel sheet (thickness 3.2 mm). At this time, the finish rolling temperature was set at about 930 DEG C, which is a straight line of Ar3. Thereafter, hot rolled steel sheets were rolled, cold rolled, continuous annealed, cooled and temper rolled under the conditions shown in Table 2 below to produce thin steel sheets.

Then, the number and distribution of precipitates and the texture of aggregates were observed and measured for each of the manufactured steel sheets, and the results are shown in Table 3 below. More specifically, the ratio of the number of carbides and the number of FeTiP precipitates were determined by counting the number of precipitates per unit length (μm) after observing the precipitates in the replica using a TEM, and calculating the average value of the number of precipitates per unit length (μm) The intensity ratio (ODF utilization) of each bearing was calculated and analyzed using EBSD based on ND direction decision bearing under the conditions of R (Rolling), T (Transverse) and N (Vertical) at 4t point. Meanwhile, FIG. 1 is a graph analyzing the degree of development of the collective organization of Inventive Example 1, and all of the inventive inventions have a tendency similar to Inventive Example 1.

Then, the r value and the bake hardening amount (BH) were measured for each of the manufactured steel sheets. The r value was measured using ASTM STD specimen, and the hardening amount was 2% pre-strain and then the yield strength value and the specimen were held again at 170 ° C. for 20 minutes The difference in yield strength values was evaluated by the difference.

Steel grade Alloy composition (% by weight) ① ^* C Si Mn P S N Sol.Al Ti Nb B Inventive Steel 1 0.0021 0.023 0.13 0.08 0.0040 0.0018 0.032 0.028 0.01 0.0008 -0.0094 Invention river 2 0.0023 0.016 0.21 0.035 0.0058 0.0028 0.032 0.02 0.011 - 0.0017 Invention steel 3 0.0032 0.012 0.22 0.03 0.0053 0.0029 0.032 0.021 - 0.0006 0.0031 Inventive Steel 4 0.0019 0.016 0.62 0.032 0.0062 0.0027 0.033 0.018 0.023 0.0009 -0.0006 Invention steel 5 0.0017 0.013 0.61 0.02 0.0061 0.0032 0.036 0.029 0.022 0.0009 0.0089 Invention steel 6 0.0008 0.008 0.9 0.03 0.0036 0.0031 0.044 0.021 0.013 0.0006 0.0050 Invention steel 7 0.0015 0.01 0.31 0.05 0.0055 0.0020 0.041 0.025 - 0.0006 0.0099 Comparative River 1 0.0035 0.02 0.8 0.06 0.0045 0.0042 0.035 0.061 - - 0.0399 Comparative River 2 0.0044 0.06 0.7 0.07 0.0038 0.0038 0.024 0.05 - - 0.0313 * ① means [Ti] - (24/7) [N] - (3/2) [S]

Steel grade Coiling Cold rolling Annealing Cooling Temper rolling Remarks Temperature (℃) Average cooling rate
(° C / s) Reduction rate (%) Heating rate
(° C / sec) Temperature
(° C) time
(s) Cooling rate
(° C / sec) Reduction rate
(%) Inventive Steel 1 680 45 78.5 3.2 848 42 3.3 0.6 Inventory 1 440 52 79.2 6.2 850 45 4.6 0.6 Comparative Example 1 Invention river 2 630 65 73.5 4.3 820 23 6.4 0.8 Comparative Example 2 630 59 79.8 3.1 861 78 4.5 0.8 Inventory 2 Invention steel 3 620 85 73.3 3.3 795 65 3.8 1.7 Comparative Example 3 680 125 79.6 2.8 849 56 3.6 0.5 Inventory 3 Inventive Steel 4 700 123 80.0 2.8 862 58 5.0 0.8 Honorable 4 765 8 80.1 5.2 845 25 1.5 0.9 Comparative Example 4 Invention steel 5 630 79 80.4 1.8 852 58 4.5 1.2 Inventory 5 720 220 80.3 5.3 810 52 3.6 0.8 Comparative Example 5 Invention steel 6 630 162 80.3 6.3 851 62 1.7 0.7 Comparative Example 6 620 140 80.2 3.6 843 55 5.0 0.6 Inventory 6 Invention steel 7 600 98 78.5 4.2 845 86 12.2 0.2 Comparative Example 7 640 102 79.3 2.8 845 51 3.0 0.9 Honorable 7 Comparative River 1 630 78 76.3 3.5 832 55 5.0 0.8 Comparative Example 8 720 75 77.8 3.5 835 56 5.0 0.8 Comparative Example 9 Comparative River 2 560 79 78.2 4.5 842 58 5.0 0.8 Comparative Example 10 560 82 78.3 4.5 832 57 5.0 0.8 Comparative Example 11

Steel grade P (%) Number of FeTiP precipitates
(Pieces / μm ² ) random
Intensity ratio
(b / a) ^* YS x r value BH (MPa) R value Remarks Inventive Steel 1 85.5 0.16 3.5 382 6.5 1.95 Inventory 1 75.3 0.31 1.6 356 4.5 1.68 Comparative Example 1 Invention river 2 91.2 0.02 1.3 315 0.3 1.65 Comparative Example 2 91.3 0.02 4.2 345 11.5 2.06 Inventory 2 Invention steel 3 86.2 0.03 1.7 286 0 1.71 Comparative Example 3 88.5 0.05 2.8 315 8.5 2.05 Inventory 3 Inventive Steel 4 86.3 0.05 3.5 336 15.7 2.23 Honorable 4 83.2 0.02 2.7 318 2.5 1.81 Comparative Example 4 Invention steel 5 89.2 0.06 3.1 326 5.2 2.23 Inventory 5 82.5 0.05 1.6 318 1.6 1.72 Comparative Example 5 Invention steel 6 81.3 0.08 3.3 284 3.1 1.92 Comparative Example 6 82.6 0.07 8.5 315 15.2 2.16 Inventory 6 Invention steel 7 85.4 0.18 6.3 312 3.2 2.05 Comparative Example 7 83.2 0.16 7.1 328 13.5 2.22 Honorable 7 Comparative River 1 63.6 8.3 3.6 325 0 1.78 Comparative Example 8 68.2 7.8 3.2 326 0 1.81 Comparative Example 9 Comparative River 2 73.5 11.5 3.6 316 0 1.85 Comparative Example 10 71.5 1.7 1.3 312 0 1.79 Comparative Example 11 * Random intensity ratio (b / a) is the average random intensity ratio of the (001) [1-10] to (110) [1-10] bearing group at t / 4 (t: thickness of thin steel plate) (b / a) of the average random intensity ratio (b) of the (111) [1-10] to (111) [- 1-12]

When referring to Table 3, the ratio and the average of the alloy composition and the carbides having a size of less than 20nm for the case of Examples 1 to 7 satisfying the production conditions units FeTiP precipitate the number per unit area, present in the ferrite grain, offering to the invention The r-value can be secured to 1.9 or more (the yield strength * r value) can be secured not less than 290 MPa, and the BH hardness can be secured to 4 MPa Or more.

However, in the case of Comparative Examples 1 to 7, the alloy composition satisfies the range suggested in the present invention, but at least one of the manufacturing conditions does not satisfy the range suggested by the present invention, and the drawability and the hardening of the hardening are poor . In addition, in the case of Comparative Examples 8 to 11, the alloy composition did not satisfy the range suggested by the present invention, so that the drawability and the hardenability of the hardening were poor.

Claims (12)

0.001 to 0.004% of C, 0.5% or less of Si (excluding 0%), 1.2% or less of Mn (excluding 0%), P of 0.005 to 0.12% , 0.1% or less (excluding 0%) of Al-soluble Al, 0.01-0.04% of Ti, the balance Fe and unavoidable impurities,
The content of Ti, N and S satisfies the following relational expression 1,
(111) [1] to the average random intensity ratio (a) of the (001) 1-10 1-10 (1-10) bearing group at the point of t / 4 (B / a) of the average random intensity ratio (b / a) of the curable resin (C) -10 to the (111)
A high strength thin steel sheet containing FeTiP precipitates of 0.2 / μm ² or less.
[Relation 1] -0.02? [Ti] - (24/7) [N] - (3/2) [S]? 0.025
(Where each of [Ti], [N] and [S] means the content (weight%) of the corresponding element)
The method according to claim 1,
0.005 to 0.04% of Nb, and 0.002% or less of B (exclusive of 0%) in terms of% by weight.
The method according to claim 1,
A high strength thin steel sheet having P of 80% or more as defined by the following formula (1).
P _in (%) = {N _in / (N _in + N _gb )} x 100
(Where N _in is the number of carbides having a circle equivalent diameter of 20 nm or less and N _gb is the number of carbides having a circle equivalent diameter of 20 nm or less present in grain boundaries)
delete The method according to claim 1,
A high-strength thin steel plate having a Yield Strength (YS) and an Average Plastic Anisotropy Coefficient (Lankford value, r-value) of 290 MPa or more.
0.001 to 0.004% of C, 0.5% or less of Si (excluding 0%), 1.2% or less of Mn (excluding 0%), P of 0.005 to 0.12% Hot-rolling a steel slab containing 0.1% or less of an acid soluble Al (exclusive of 0%), 0.01-0.04% of Ti, the balance Fe and unavoidable impurities to obtain a hot-rolled steel sheet;
Winding the hot-rolled steel sheet at a temperature of 450 to 750 ° C;
Cold rolling the rolled hot-rolled steel sheet at a reduction ratio of 75% or more to obtain a cold-rolled steel sheet;
Heating the cold-rolled steel sheet to an annealing temperature of 830 to 880 캜, and continuing annealing at the annealing temperature for an annealing time of 30 to 80 seconds;
Cooling the continuously annealed cold rolled steel sheet to a temperature of 650 ° C or lower at a rate of 2 to 10 ° C / sec;
And subjecting the cooled cold-rolled steel sheet to rough rolling at a reduction ratio of 0.3 to 1.6% without plating,
Wherein the average heating rate from the (recrystallization starting temperature + 20) ° C to the annealing temperature is 5 ° C / sec or less at the time of increasing the temperature of the cold-rolled steel sheet.
The method according to claim 6,
By weight, at least one selected from the group consisting of Nb: 0.005 to 0.04% and B: 0.002% or less (excluding 0%).
The method according to claim 6,
Wherein the finish rolling temperature during the hot rolling is Ar 3 ° C or higher.
9. The method of claim 8,
And the average cooling rate from the finish rolling temperature to the coiling temperature is 10-200 占 폚 / sec.
The method according to claim 6,
Wherein the annealing temperature (T, 占 폚) and the annealing time (t, sec) at the time of the continuous annealing satisfy the following relational expression (2).
[Relation 2] 30? 0.001 * T * t? 60
The method according to claim 6,
Further comprising a step of hot-dip galvanizing the surface of the rough-rolled cold-rolled steel sheet.
12. The method of claim 11,
Further comprising the step of performing an alloying heat treatment at 450 to 600 占 폚 after the hot dip galvanizing.

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